【作者】 孙晶；

【导师】 刘正；

【作者基本信息】 沈阳工业大学 ， 材料加工工程， 2011， 博士

【摘要】 Mg-Al系合金是应用最早、也是迄今为止牌号最多、应用最为广泛的一类镁合金,但这类镁合金结晶凝固温度范围较宽,晶粒有粗化的倾向,凝固过程中粗大的β-Mg17Al12相以离异共晶的方式沿晶界呈网状析出,导致合金的力学性能变差,极大地限制了Mg-Al系合金作为结构件在汽车、飞机等领域的应用。最常用的提高镁合金力学性能的方法是合金化。在常用的合金元素中,能同时提高镁合金强度和塑性的有Al、Zn、Ca、Ag、Ce、Ga、Ni、Cu、Th等,目前有关Al、Zn、Ca、Ag、Ce、Ni、Cu、Th对镁合金组织及性能影响的研究较多,而Ga对镁合金组织及性能影响的研究则较少。为此,本文选择应用最广泛的Mg-Al系合金,研究Ga对其微观组织及力学性能的影响规律。试验配制了Mg-8%Al-(0-4)%Ga、Mg-(8-12)%Al-2%Ga及AZ91-(0-2)%Ga镁合金,分析了Ga对Mg-Al系合金铸态微观组织及力学性能的影响规律,在此基础上重点研究了热处理工艺及轧制处理对Mg-8%Al-2%Ga合金第二相析出规律的影响,探讨了其强化机制。研究结果表明,Mg-Al合金的铸态组织由α固溶体、沿晶界分布的以β-Mg17Al12为基的共晶组织及共晶组织附近的次生相所组成。次生相表现为与α固溶体呈片层相间的形式;共晶组织则是由于快冷先形成离异的β-Mg17Al12相,转而形成以β-Mg17Al12相为基的共晶组织。加入的Ga元素主要存在于β相中,以及β相与α相交界处,可以明显地细化晶粒,改善Mg-Al合金的铸态组织,同时抑制共晶组织的形成,使离异共晶的β-Mg17Al12相趋于细小、光滑;对于Mg-8%Al-(0-4)%Ga合金而言,随着Ga含量的增多,Ga将向β相与基体α相交界处富集,并当含量达到4wt.%时,β相发生熔断,形成Ga2Mg新相,并在熔断处和β相与α相交界处可观察到Mg-Al-Ga三相共晶体。对于Mg-8%Al合金而言,Ga的最佳加入量为2wt.%。Mg-(8-12)%Al合金中加入2wt.%Ga后,可抑制共晶组织的形成,使Mg-Al合金的室温力学性能明显提高。此时,铸态下合金的抗拉强度σb为248.4MPa、伸长率δ为10.4%,T6状态下的抗拉强度σb为288.0MPa、伸长率δ为4.7%。Mg-8%Al-2%Ga合金固溶处理时,首先回溶的是铸态组织中次生的片状β-Mg17Al12相,然后是离异共晶β-Mg17Al12相;并且随着固溶时间的延长,晶粒逐渐长大,在随后的时效处理时以非连续方式析出的β-Mg17Al12相逐渐减少。Mg-8%Al-2%Ga合金经415℃固溶处理后,以不同速度冷却时将得到不同的组织:加热后水冷,得到的是单一的过饱和固溶体,晶界呈平直状;加热后空冷,晶界处有非连续的β-Mg17Al12析出,晶界呈波纹状起伏;加热后炉冷时,晶内有大量片层状β-Mg17Al12相析出,形态类似于珠光体组织。Mg-8%Al-2%Ga合金轧制变形后,组织中出现大量的孪晶,表明轧制时变形方式以孪生为主。轧制变形改变了Mg-8%Al-2%Ga合金时效时第二相的析出位置及形态,未变形时β相析出方式为非连续析出和连续析出,轧制变形后析出方式以在缺陷处(晶界、孪晶界、孪晶)析出细小、弥散的β相为主。轧制变形可以有效地提高Mg-8%Al-2%Ga合金的硬度。Mg-8%Al-2%Ga合金轧制变形后在150℃-300℃退火时,未发生再结晶;400℃退火时,发生了再结晶。150℃-200℃退火时,仅在晶界及孪晶处析出细小、弥散的β相;300℃退火时,除了晶界、孪晶析出细小、弥散的β相外,在晶粒内部沿基体α相的(0001)晶面析出长条状β相,退火3h组织中还出现了一些菱形片状彼此夹角为120°(60°)的β相。当合金于150℃、200℃、300℃退火时,硬度随退火时间的延长先下降,然后上升,到一定时间达到峰值,然后基本不变。这三个温度退火硬度的变化规律是退火温度越高,硬度值也越高。这是因为退火温度高,析出的β相数量增加,导致合金的硬度值提高。400℃退火时,发生了再结晶,使硬度值明显低于前三个温度。对于AZ91镁合金而言,加入2wt.%Ga后可以明显地改变其铸态组织,共晶组织消失,离异共晶的β相为光滑的块状并有球化倾向,在β相周围出现了片层状次生相,加入的Ga主要存在于β相的周围,还有一些存在于β相中;Ga的加入,使Zn元素向β相的周围偏聚;2wt.%Ga可以明显提高AZ91镁合金的室温力学性能。更多还原

【Abstract】 The Mg-Al series magnesium alloys are the most extensively used magnesium alloy so far. However, because of their widen solidification range, theβ-Mg17Al12 phase frequently ll precipitate along theα-Mg grain boundary in the shape of network by the reaction of divorce eutectic, result in the poor mechanical properties of the Mg-Al series magnesium alloys component. This limited the Mg-Al series magnesium alloys application in the industry of automobile and aerospace. At present, the most frequently used method to increase the mechanical properties of magnesium alloy is alloying. During the used alloying elements, those who can increase the strength and ductility of magnesium alloys are Al、Zn、Ca、Ag、Ce、Ga、Ni、Cu、Th, ect. There are lots of literatures have studied the effects of elements of Al、Zn、Ca、Ag、Ce、Ni、Cu、Th on the microstructure and mechanical properties of Mg–Al series magnesium alloys, but about the effect of Ga on the microstructure and mechanical properties of magnesium alloys are very limited. The aim of the present thesis is to investigate the effect of Ga on the microstructure and mechanical properties of Mg–Al series magnesium alloys.The nominal chemical composition of magnesium alloys studied in this thesis are Mg-8%Al-(0-4)%Ga、Mg-(8-12)%Al-2%Ga and AZ91-(0-2)%Ga. The thesis analysis the effect of Ga on the microstructure and mechanical properties of casting magnesium alloys, the emphasis was laid on the effect of heat treatment and rolling technology on the second phase precipitation pattern of Mg-8%Al-2%Ga magnesium alloy, and the strengthening mechanism of the alloys was discussed at the same time. The experimental results demonstrated that the microstructure of casting Mg-Al magnesium alloy is consisted ofα-Mg, the eutectic ofβ-Mg17Al12 andα-Mg, and the secondary phase adjacent of eutectic phase. The secondary phase is in the shape of lamella alternate withα-Mg; the eutectic is transformed by the primary divorce eutecticβ-Mg17Al12 phase which is formed under the condition of high solidification rate.Most of the Ga addition is mainly concentrated inβphase and in the interface ofαandβphase. Ga addition not only can refine the casing microstructure but also can suppress the formation of eutectic and decrease the grain size of divorced eutectic phase ofβ-Mg17Al12. For the Mg-8%Al-(0-4)%Ga magnesium alloy, with increasing the Ga addition, it is mainly rich in the interface ofβphase andαmatrix. When the Ga addition is increased to 4wt.%βphase is broken and Mg-Al-Ga ternary eutectic phase is observed in the interface ofαandβphase, there is a new phase of Ga2Mg formed in the alloy when the Ga addition is increased to 4wt.%.For the magnesium alloy of Mg-8%Al the optimal addition of Ga is 2wt.%. When 2wt.%Ga is added to the Mg-(8-12)%Al alloy the supersaturation of Al in the matrix is decreased, as a result the ambient mechanical properties of the alloy is increased obviously. The ultimate tensile strength and elongation of the alloy in the as-cast state is 248.4MPa and 10.4%, respectively, while the ultimate tensile strength and elongation of the alloy in the T6 state is 288.0MPa and 4.7%, respectively.When the Mg-8%Al-2%Ga magnesium alloy is undergone solid solute treatment, the first dissolved phase is second lamellaβ-Mg17Al12 phase, then is the divorced eutecticβ-Mg17Al12 phase. With increasing the solid solution time the grain size of the matrix is increasing and during the subsequent aging the amount of discontinues precipitationβ-Mg17Al12 phase is decreasing. The different microstructure will be obtained after the Mg-8%Al-2%Ga magnesium alloy is solid solution-treated at 415℃then cooling at different method, the supersaturation microstructure with the straight and smooth grain boundary will be obtained after quench cooling, while microstructure featured of undulant grain boundary with discontinues precipitation ofβ-Mg17Al12 phase after air cooling, when the alloy is furnace cooling after solid solution, there is a great amount of lamellarβ-Mg17Al12 phase precipitated within the grain.After rolling deformation the microstructure of Mg-8%Al-2%Ga magnesium alloy is consist of great amount of twins, implying that the rolling deformation is mainly by twinning. Rolling deformation change precipitation site and morphology of the secondary phase. The precipitation patterns ofβphase are continuous and discontinuous before rolling, but theβphase is mainly precipitated in the site of defects(including grain boundary, twin grain boundary and with the twins).Rolling deformation can increase the hardness of Mg-8%Al-2%Ga magnesium effectively. The recrystallization of the alloy didn’t not occure when it is annealing at 150℃to 300℃after rolling deformation, while annealing at 400℃the recrystallization take place. When the alloy is annealing at 150℃to 200℃the phase is only precipitated at grain and twin boundaries at small and dispersed morphology, when the annealing temperature is increasing to 300℃except above motionedβphase, there is, the long lamellarβphase precipitated along (0001) plane of theαmatrix is also observed, when the annealing time increased to 3 hours theβphase is in the shape of lamellar diamond and the angle between two lamellar is 120°(60°). When the alloy is annealing at 150℃、200℃、300℃, respectively, the hardness of the alloy is decreasing at first then increasing with increasing the annealing time, then lever off. The higher the annealing temperature, the higher the hardness value, that is because the higher the annealing temperature, the more theβprecipitation amount.The 2wt.%Ga addition to the AZ91 magnesium alloy can alternate the casting microstructure apparently, the eutectic is diminished and the divorced eutecticβphase tends to be block and spheroidization, and there is lamellar secondary phase formed around theβphase. The Ga is mainly concentrated aroundβphase, and Zn segregate aroundβphase because of the addition of Ga. The ambient mechanical properties of AZ91 magnesium alloy is increased apparently by 2wt.%Ga.更多还原